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GeSn alloys are promising materials for CMOS-compatible mid-infrared lasers manufacturing. Indeed, Sn alloying and tensile strain can transform them into direct bandgap semiconductors. This growing laser technology however suffers from a number of limitations, such as poor optical confinement, lack of strain, thermal, and defects management, all of which are poorly discussed in the literature. Herein, a specific GeSn-on-insulator (GeSnOI) stack using stressor layers as dielectric optical claddings is demonstrated to be suitable for a monolithically integration of planar Group-IV semiconductor lasers on a versatile photonic platform for the near- and mid-infrared spectral range. Microdisk-shape resonators on mesa structures were fabricated from GeSnOI, after bonding a Ge0.9Sn0.1 alloy layer grown on a Ge strain-relaxed-buffer, itself on a Si(001) substrate. The GeSnOI microdisk mesas exhibited significantly improved optical gain as compared to that of conventional suspended microdisk resonators formed from the as-grown layer. We further show enhanced vertical out-coupling of the disk whispering gallery mode in-plane radiation, with up to 30% vertical out-coupling efficiency. As a result, the GeSnOI approach can be a valuable asset in the development of silicon-based mid-infrared photonics that combine integrated sources in a photonic platform with complex lightwave engineering.

The holy grail of group-IV light emission would be electrically pumped lasing at Room Temperature (RT). In CEA-LETI, the performance of p-i-n light emitting diodes is limited by the use of in-situ doped Ge layers beneath and above the active layer stack [1]. To overcome Sn segregation during the growth of the top electrode and, in thick GeSn/Ge stacks, avoid plastic relaxation, it would be ideal to switch over to in-situ doped GeSn grown at temperatures similar to that of the active region. We have therefore explored the in-situ boron and phosphorous doping of GeSn on Ge strain relaxed buffers with Ge2H6+SnCl4+ B2H6 and Ge2H6+SnCl4+PH3 chemistries. Growth pressure, temperature, F(Ge2H6)/F(H2) and F(SnCl4)/F(H2) Mass Flow Ratio (MFR) were 100 Torr, 349 °C, 7.92x10-4 and 4.78x10-5. X-Ray Diffraction (XRD) measurements were conducted to gain access to the GeSn composition and layer thickness, as shown in Fig. 1 and Fig. 2. As the GeSn peak position was determined by the Sn content and also the B or P substitutional concentrations, the actual layer composition could not be unambiguously determined. Therefore, the Sn content in the following is an apparent Sn content. Thickness fringes outline the high GeSn crystalline quality. As the F(B2H6)/F(Ge2H6) or the F(PH3)/2 F(Ge2H6) MFRs increased, the apparent Sn content stayed at first roughly around 6.5 %, as shown in Fig. 3. This value was the same for i-GeSn, GeSn:B and GeSn:P, in contrast to previous literature findings [2]. As the MFRs reached the highest values, the apparent Sn content reduced significantly for GeSn:B (from 6.5 % down to 4.6 %) and less significantly GeSn:P (from 6.6 % down to 5.6 %). The significantly smaller size of B compared to P (aB = 3.852 Å aP = 5.014 Å) reduced the compressive strain in GeSn:B 2.8 times more than in GeSn:P for the same dopant concentrations, explaining the more pronounced effect. We quantified the substitutional B and P incorporation by assuming that the Sn content was not impacted by B or P incorporation. The estimated B and P substitutional concentrations then increased from 2.8x1019 cm-3 up to 3.8x1020 cm-3 and from 2.5x1019 cm-3 up to 5.4x1020 cm-3 (highest four MFRs for GeSn:B and GeSn:P), shown in Fig. 4. Those values are close indeed to the highest substitutional B or P concentrations obtained in Ge:B or Ge:P layers grown at 350°C, 100 Torr with Ge2H6: 4.8x1020 cm-3 [3] and 4.1x1020 cm-3 [4], respectively. We determined elemental growth rates by multiplying the GR by the Ge or apparent Sn content to have a more detailed look at the growth mechanics (Fig. 5). The introduction of SnCl4 catalyzed the growth rate of GeSn compared to that of pure Ge, in line with the literature [4]. As the F(B2H6)/F(Ge2H6) or the F(PH3)/2F(Ge2H6) MFR increased, the Sn GR component stayed constant at around 2.5 nm min.-1. Meanwhile, the Ge GR component was constant at around 37 nm min.-1 for lower MFRs. For high dopant flows, the Ge GR component increased significantly for GeSn:B and reached 48 nm min.-1. The increase was less significant for GeSn:P, with at most 40 nm min.-1. This is to our knowledge, the first time an increase of the GeSn:P GR is reported for high dopant flows. Similar increases were found for Ge:P and Ge:B grown at 350°C, however [3-4]. The constant Sn GR component and the increase of the Ge GR component at high MFRs are likely due to an increased incorporation of Ge thanks to B and P opening surface sites (Indeed, the Ge GR component is surface rate limited, while the Sn GR component is reaction rate limited [5]). A cross-hatch along was found by Atomic Force Microscopy for all layers, which were otherwise smooth and featureless, as shown in Fig. 6 top for GeSn:B and Fig. 6 bottom for GeSn:P. The only exceptions were the GeSn:B samples grown at the three highest MFRs. There were then some islands, which were likely due to B and/or Sn surface segregation. Other studies will be performed to gain access to the Sn, P and B contents. [1] M. Bertrand et al., 2018 IEEE International Conference on Group-IV Photonics, pp. 45-46. [2] J. Margetis et al., Mater. Sci. Semicond. Process. 70, 38 (2017). [3] J.M. Hartmann et al., ECS PRIME 2020 Meeting Abstract ID 138196, Symposium G03. [4] J. Aubin et al., ECS J. Solid State Sci. Technol. 6, 1 (2017). [5] J. Margetis et al., J. Vac. Sci. Technol. A 37, 021508 (2019). Figure 1

International audience; Tensile strained GeSn lasers could be used as tunable infrared light sources in silicon photonics. We explore here the relationship between the uniaxial tensile strain applied to suspended microbridges and the maximum lasing temperature reached in GeSn 13% cavities.

Direct-band gap GeSn alloys have strong potentials as MIR light emitters. We investigate LEDs with 13% and 16% of tin. Light emission is observed up to room temperature for both types of LEDs. The Ge 0.84 Sn 0.16 LED is used in a gas sensor to detect CH4.

In-situ boron doping of Si and Ge can be of use in various types of devices. Poly-Si:B gates can be integrated in p-type metal oxide semiconductor (pMOS) transistors. Monocrystalline Si:B layers can be used in p-i-n junctions for RF switches, attenuators, photo-detectors (PDs) or phase shifters. Ge:B layers can be integrated in the sources and drains of Ge pMOS transistors or as the p-type films in Ge p-i-n PDs. If the growth temperature is low enough, they can even be used in GeSn PDs or light emitting devices. We have explored here the boron-doping of Si and Ge with Si2H6 + B2H6 and Ge2H6 + B2H6. Such chemistries enable Si (Ge) epitaxy at temperatures lower than with conventional SiH4 (GeH4) precursors. Growth temperatures and pressures in our 200mm RP-CVD tool were 550°C, 20 Torr (Si) and 350°C, 100 Torr (Ge). X-Ray Diffraction (XRD) was used to convert the tensile strain in Si:B or Ge:B into substitutional B concentrations, as shown in Fig. 1 and Fig. 2, with layer peaks moving away from the Si(001) substrate or the Ge Strain-Relaxed Buffers (SRB) as the diborane flow increased. The presence of marked thickness fringes testified to the high crystalline quality of the films. Secondary Ions Mass Spectrometry (SIMS) gave us the atomic B concentration. X-Ray Reflectivity on Si0.5Ge0.5 / Si:B stacks yielded Si:B Growth Rates (GR). Finally, differential weighting gave us the thickness of Ge:B layers grown on Ge SRBs and other GR values. A huge Si:B GR increase, from 9.5 up to 40 nm min.-1, was evidenced when adding relatively small amounts of B2H6 to Si2H6, as shown in Fig. 3. This was due to the presence of B atoms on the surface which catalyzed H desorption, freeing sites for growth. Ultra-high substitutional and atomic boron contents were obtained in those Si:B layers (at most: 2.7x1020 and 1.1x1021 atoms cm-3), as shown in Fig. 4. Films became poly-cristalline for the highest F(B2H6)/F(Si2H6) Mass-Flow Ratios (MFRs) probed, without XRD peaks anymore and a boron ion concentration which was reduced compared to its peak value of 2x1020 cm-3. The good agreement between substitutional, atomic and ionic B concentrations for low to medium MFRs is to be underlined. By contrast, the growing discrepancy between XRD and SIMS for the highest diborane flows probed might partly be due to a SIMS artefact. Indeed, the Relative Sensistivity Factor used to convert the raw SIMS signals into B concentrations was for doped layers and not higher boron content binary alloys. The presence of B atoms in nanoclusters that did not contribute to XRD cannot be excluded, however. A definite growth rate increase (from 6 up to 21 nm min.-1) and really high substitutional B concentrations (at most: 4.8x1020 atoms cm-3) were also obtained for Ge:B, as shown in Fig. 5 and Fig. 6. The good agreement between XRD and SIMS for medium F(B2H6)/F(Ge2H6) MFR is to be underlined. The growing discrepancy for high diborane flows might be due to the same SIMS artifact as before and to the presence of B atoms in nanoclusters. x/(1-x) = m*F(B2H6)/F(Si2H6 or Ge2H6) relationships accounted for the semi-linear increase of the substitutional or atomic B concentration x with the diborane flow, with (i) msubst. = 1.85 (matom. = 7.45) in Si:B and (ii) msubst. = 0.15 (matom. = 0.68) in Ge:B. m values above 1 meant that B incorporation was easier than Si incorporation in Si:B. In contrast, B incorporation was more difficult than Ge incorporation in Ge:B alloys, as m was less than 1, then. Current growth rate increases and B concentrations are definitely higher than with mainstream chemistries. For Si:B grown at 650°C, 20 Torr with SiH4 + B2H6, we had for instance GR increases from 9 up to 23 nm min.-1, with at most 9x1019 cm-3 B ions concentrations before layers became poly-crystalline (ECS Trans. 16 (10), 485 (2008)). For Ge:B grown at 400°C, 100 Torr with SiH4 + B2H6, GR increased from 10 up to 18 nm min.-1, with at most 1020 cm-3 B ions concentrations (Thin Solid Films 557, 4 (2014)). The advantages of switching over to higher order Si and Ge precursors and adopting lower growth temperatures are obvious. Figure 1